Process for Production of a Carboxylic Acid/Diol Mixture Suitable for Use in Polyester Production

ABSTRACT

Disclosed is a wrought magnesium alloy having excellent strength and extrusion or rolling formability, and a method of producing the same. The wrought magnesium alloy comprises 0.1-1.5 at % group IIIa, 1.0-4.0 at % group IIIb, 0.35 at % or less of one selected from the group consisting of groups IIa, IVa, VIIa, IVb, and a mixture thereof, 1.0 at % or less of group IIb, and a balance of Mg and unavoidable impurities and thus has a second phase composite microstructure. The wrought magnesium alloy of the present invention has high strength, toughness, and formability in addition to the electromagnetic wave shield ability of magnesium. Accordingly, the wrought magnesium alloy is a material useful to portable electronic goods, such as notebook personal computers, mobile phones, digital cameras, camcorders, CD players, PDA, or MP3 players, automotive parts, such as engine room hoods, oil pans, or inner panel of door, or structural parts for airplane.

TECHNICAL FIELD

The present invention relates to a wrought magnesium alloy, whichcontains a second phase consisting of an intermetallic compound, therebyhaving excellent strength, formability, and corrosion resistance. Moreparticularly, the present invention pertains to a wrought magnesiumalloy, which comprises 0.1-1.5 at % group IIIa. 1.0-4.0 at % group IIIb,0.35 at % or less of one selected from the group consisting of groupsIIa, IVa, VIIa, IVb, and a mixture thereof, 1.0 at % or less group IIb,and a balance of Mg and unavoidable impurities and which thus has asecond phase composite microstructure consisting of an intermetalliccompound, and a method of producing the same.

BACKGROUND ART

Having a density of 1.74-1.95 g/cm³ or so, a relatively low specificgravity that is ⅔ of that of aluminum, excellent specific strength andmachinability, magnesium alloys have been developed as light structuralmaterials for airplanes and automobiles. However, since magnesium has ahexagonal close packed (HCP) lattice crystal structure, formability isvery low, and thus, the application is limited to a field in which theforming is achieved using a casting process. Particularly, its practicaluse is limited due to problems, such as severe oxidation of melts,reduction of strength at high temperatures, and low corrosionresistance. Effort has been made to avoid the above disadvantages, thusenabling stable dissolution in atmospheric air, while employing a sulfurhexafluoride (SF₆) gas, a carbon dioxide gas, an argon gas and the like,and production of a plate by Direct Chilled casting process.

Of the magnesium alloys, a Mg—Zn alloy shows an excellent age hardeningbehavior, and is advantageous in that since a microstructure is refinedthrough heat treatments, strength and ductility significantly increaseand it is easy to work and weld. On the other hand, it isdisadvantageous in that since micropores are formed as a casting processdue to the addition of Zn, it is difficult to apply the Mg—Zn alloy to acasting process, such as die-casting. Additionally, it is difficult todesirably improve strength because it is grown as the coarse grain. Toovercome the above disadvantages, studies have been made to improveformability using a grain boundary slip, in which some alloy elementsare added to Mg—Zn binary alloys to refine grains. With respect to this,J. P. Doan and G. Ansel suggest a method of improving the strength of analloy, in which Zr is added to refine grains constituting a Mg—Zn alloy(J. P. Doan and G. Ansel, Trans, AIME, vol. 171 (1947), pp. 286-295).However, since Zr has a high melting point and a low solubility to Mg atroom temperature, it mostly exists at a grain boundary, thus acting as afracture initiation site when external stress is applied. In thisregard, it becomes possible to conduct a plastic working process usingthe ductility of a single-phase solid solution after alloy materials,including aluminum and zinc or manganese, such as AZ31B or AM20, aredeveloped. However, even though their microstructures have asingle-phase solid solution and thus have excellent ductility, they aredisadvantageous in that since a strain hardening ability is poor and itis difficult to prevent the grain growth, formability is poor due toanisotropy. A technology, in which different portions are heated atdifferent temperatures to achieve a warm working process, is suggestedso as to avoid the above disadvantage. However, the technology isproblematic in that the different heating temperatures of the differentportions significantly increase the production cost of a press mold. Asan alternative method, a thixo-molding method, in which preliminarilyflake shaped powder is compacted at high temperatures in a region whereliquid and solid phases coexist, is suggested. However, this method isdisadvantageous in that the powder is expensive, and in that it isdifficult to apply an electroplating process because the powder pressedmaterial has a porous structure. Magnesium has low corrosion resistance,and thus, it is necessary to treat a surface of magnesium, butundesirably, a gas phase plating process or an electroless platingprocess requires chemical and treatment costs that are much higher thanthe electroplating process. However, products having high porosity andlow density, such as die-casting and thixo-molding products, aredifficult to apply to a wetplating process because corrosion occurs dueto chemicals soaked into the pores.

Furthermore, Korean Pat. Laid-Open Publication No. 2003-0048412discloses an alloy, which contains 3.0-10.0 wt % Zn, 0.25-3.0 wt % Mn,Al, Si, and Ca. However, even though the alloy containing Zn in anamount of 2% or more has high strength, it has a disadvantage in thatfree zinc (Zn) readily forms a low melting point eutectic phase. Forexample, if Mg₇Zn₃, having a low melting point that is less than 350°C., exists, corrosion resistance is low. And the plate is easily crackedat both sides during a rough rolling process for breaking the coarsedendrite structure, so draw-ability is poor because of high anisotropy.Korean Pat. Laid-Open Publication No. 2002-0078936 (U.S. Pat. No.6,471,797) discloses a method of improving strength and formabilityusing a Mg—Zn—Y eutectic ternary alloy quasi-crystalline phase whichcontains 1-10 at % Zn and 0.1-3 at % Y. However, this method isdisadvantageous in that the amount of Zn must be enough so as todesirably assure a quasi-crystalline phase effect. The composition of acast product is not uniform because a specific gravity between zinc andmagnesium is significantly different. The micro-pores at the grainboundary reduce corrosion resistance, and tears form at sides of a plateduring the hot rolling process. On the other hand, in Korean Pat.Application No. 10-2003-0044997 which has been made by the inventor ofthe present invention, tears caused by an ununiform composition duringrolling process is reduced by reducing the amount of Zn. However, sincethe second phase improving plasticity is a low melting point eutecticphase that is formed at grain boundary after the matrix is formed, thesecond phases are dispersed through the break down rolling aftersolidification. Accordingly, it is difficult to uniformly disperse them.Hence, the initial rolling process must be repeated a few times within areduction ratio of 5-10%, and the process must be conducted within areduction ratio range of 15-20% after the cast structure is broken so asto get a good quality without side cracks.

Many other patents disclose a method of producing a light magnesiumalloy strip or powder having high strength, in which an amorphousstructure is produced through the Rapid Solidification process. KoreanPat. Laid-Open Publication No. 1990-0004953, entitled ‘high strengthmagnesium alloy’, Korean Pat. Laid-Open Publication No. 1993-846,entitled ‘magnesium alloy having high strength’, Japanese Pat. Laid-OpenPublication No. H05-70880, entitled ‘magnesium alloy material havinghigh strength and method of producing the same’, Japanese Pat. Laid-OpenPublication No. H06-41701, entitled ‘amorphous magnesium alloy havinghigh strength and method of producing the same’, Japanese Pat. Laid-OpenPublication No. H07-54026, entitled magnesium alloy having high strengthand method of producing the same. U.S. Pat. Nos. 4,675,157, 4,765,954,4,853,035, 4,857,109, 4,938,809, 5,071,474, 5,078,806, 5,078,807,5,087,304, 5,129,960, and 5,316,598, EP No. 0,361,136A1, and French Pat.No. 2,688,233 disclose the formation of an amorphous structure through aRapid Solidification process. Since the cooling rate must be conductedat 10⁵-10⁷° C./s to form the amorphous structure, the patents are usefulto produce powder or a thin strip, but not to produce a common plateshape. Accordingly, an ingot, which is produced by compacting amorphouspowder under the recrystallization temperature, is employed in order toconduct a rolling or a press forming.

Furthermore, U.S. Pat. Nos. 637,040, 3,391,034, 4,116,731, 4,194,908,and 5,059,390, and English Pat. No. 2095288 disclose the fact that somerare-earth elements are used to prevent the grain growth or the grainboundary slip at high temperatures while their eutectic phases exist atgrain boundary so as to improve creep resistance. However, the eutecticphase mostly has a coarse microstructure that is incoherent to a matrixmicrostructure, and thus, formability is insufficiently improved. Aswell, Japanese Pat. Laid-Open Publication Nos. H7-109538A, U.S. Pat.Nos. 5,693,158, 5,800,640 and 6,395,224 disclose a method of producinggoods having low crack sensitivity, in which Sr, Li or B is employed andheat treatments are conducted to refine a particle size of crystals of acast product. However, these patents are useful to cast products, butcannot be directly applied to wrought products. Japanese Pat Laid-OpenPublication No. H10-147830A discloses the use of 6-12 wt % Y and 1-6 wt% Gd, and hot forging and subsequent aging processes to improve creepresistance to be applied to engine parts. However, the patent cannot beapplied to wrought products because the product cost significantlyincreases due to the use of a lot expensive elements, and the coarseintermetallic compounds are incoherent to the matrix. Furthermore, amethod of improving formability, in which an excessive amount of Li isemployed to change the lattice structure of a matrix microstructure intoa body-centered cubic lattice, is suggested. However, this method is notuseful to casing materials when considering a Galvanic reaction of Liand an increased cost due to the use of an excessive amount of Li.

DISCLOSURE OF INVENTION Technical Problem

Accordingly, the present invention has been made keeping in mind theabove problems occurring in the prior art, and an object of the presentinvention is to provide a wrought magnesium alloy which contains theintermetallic compound coherent to a matrix microstructure and which hasa second phase composite microstructure, thereby improving elongationand anisotropy to assure excellent formability and corrosion resistance.In order to accomplish the above object, an alloy consisting of three ormore elements is used to activate a slip plane. Additionally, in orderto activate the slip plane according to increasing in temperature, IIIaand IIIb groups are added together to reduce stacking fault energy andto improve corrosion resistance of the matrix microstructure.Furthermore, fine intermetallic compound particles dispersed duringextrusion and rolling processes are employed to improve strain hardeningability and formability.

Technical Solution

In order to accomplish the above object, the present invention providesa wrought magnesium alloy having excellent formability and platingproperties, which comprises 0.1-1.5 at % group IIIa, 1.0-4.0 at % groupIIIb, 0.35 at % or less of one selected from the group consisting ofgroups IIa, IVa, VIIa, IVb, and a mixture thereof, 1.0 at % or less ofgroup IIb, and a balance of Mg and impurities and which thus has asecond phase intermetallic compound.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 illustrates a box sample formed using a wrought magnesium alloysheet, according to the present invention;

FIG. 2 illustrates a cup-shaped sample formed using the wroughtmagnesium alloy sheet, according to the present invention;

FIG. 3 illustrates the box sample formed using AZ31 sheet;

FIG. 4 illustrates a microstructure of a material of No. 1 in Table 1,which is cast and then diffusion annealed at 400° C. for 5 hours;

FIG. 5 illustrates a microstructure of an extruded material according tothe present invention, which is annealed; and

FIG. 6 illustrates a microstructure of a rolled sheet according to thepresent invention.

BEST MODE FOR CARRYING OUT THE INVENTION

The present invention is characterized in that the fine 2nd phaseprecipitates, which is coherent to matrix microstructure, is formed inthe solid solution microstructure having excellent ductility, therebymaking grains fine and improving formability. When the grains are madefine, the strength of most materials increases. The reason is that adislocation moves along the specific slip plane in the course of plasticdeformation of the metal in such a way that the dislocation does notdirectly move from one grain to another grain. But direction ofdislocation changes its route because of the grain boundary barriereffect. Accordingly, since the grain boundaries act as barriers in themovement of the dislocation, dislocations are pile up at a grainboundary, thereby preventing deformation. The high temperature stablephase must be capable of being formed in order to make the grain fine,and desired solid solubility must be assured at high temperatures inorder to be coherent to a matrix microstructure. Furthermore, a sizedifference between elements of a matrix metal and atoms must be about15% so as to assure a desirable matrix reinforcement effect. Manystudies have been made of the effect of an intermetallic compound on asolid solution. Particularly, a matrix reinforcement effect caused bythe dispersion of fine intermetallic compound particles is well known inthe metallurgy engineering (Mechanical Metallurgy, 2nd ed., George E.Dieter, McGraw-Hill, 1981, pp. 221-227). The intermetallic compound hasa high melting point and strong bonding strength, thus having highhardness and thermally stable. Because of the finely dispersed secondphase particles, these alloys are much more resistant torecrystallization and grain growth than single-phase alloys. However, ifthe intermetallic compound has a microstructure that is incoherent tothe matrix microstructure, it acts as a fracture initiation site andthus has increased strength, but elongation or total ductility isreduced even though the matrix microstructure has ductility.

If the second phase of conventional magnesium alloy is not a highmelting point phase in the matrix microstructure. The 2^(nd) phase islow melting point eutectic phase during the solidification instead ofthe precipitates. Therefore, the eutectic phase is mostly incoherent tothe matrix microstructure. It is scarcely an atomic match for the matrixmicrostructure, thus effectively preventing grain growth or over-aging.However, disadvantageously, it reduces the formability of a material oracts as a fracture initiation site. And thus, these type alloys areunsuitable for wrought magnesium alloy. Even though a duplexmicrostructure is formed, the movement of the dislocation isineffectively prevented if the second phase is not strong, resulting inundesirably improved anisotropy or strength.

It is known that group IIIa elements employed in the present inventionreadily form intermetallic compounds having a cubic lattice and thushave high a matrix reinforcement effect and ductility. Alan Russel andKarl Gschneidner Jr. of the Ames Laboratory of Iowa State University,which is affiliated with the U.S. Department of Energy, reported that anintermetallic compound formed by the group IIIa has a B2 cubic lattice,such as CsCl, unlike a B27, B33, or DO₁₁ orthorhombic lattice of aconventional intermetallic compound. And thus, it has excellentductility (Nature Materials, 2, Sep. 2003, PP 587-590). Currently, ithas been reported that many intermetallic compounds containing group maelements are coherent to the magnesium matrix, and it is presumed thatthe ductility of the intermetallic compound is caused by stackingfaults.

Furthermore, many researchers, including A. P. Tsai, ascertained thefact that since a quasi-crystal intermetallic compound formed by groupIIIa elements has high adhesion energy and Young's modulus, it is asubstance having high strength and ductility. Based on the above fact,many studies have been made to apply the quasi-crystal intermetalliccompound to a structural material. Particularly, in the magnesium alloyfield, Japan and Korea have taken the lead in the studies of a Mg—Zn—Yalloy containing quasi-crystal particles (Materials Science andEngineering A300, 2001, pp. 312-315; Acta Materialia 50 (2002) pp.2343-2356; Materials transactions vol. 42, No. 10 (2001) pp. 2144-2147;TMS 2002 conference, Magnesium Technology 2002, pp. 141-150; Journal ofAlloys and Compounds 342 (2002) pp. 445-450).

The above studies have proven the following fact. Since zinc separationoccurs during a dissolution process due to the high Zn content (4 at %or more), the composition is ununiform. And free Zn forms a low meltingpoint eutectic phase, thus undesirably causing a side crack duringrolling process. However, Y, the group IIa elements, forms a icosahedralquasi-crystal phase in conjunction with Mg and Zn, and thus, the phasereinforces a matrix while being coherent to the matrix, therebyeffectively preventing grain growth at a high temperature until 400° C.Particularly, A. Inoue of Japan confirmed using a high resolutionelectron microscope (HREM) that in a magnesium alloy, which is producedthrough an RSP process and contains 2 at % Y and 1 at % Zn, the ABACABtype of stacking faults are formed every 6 periods (Scripta Materialia49 (2003) pp. 417-422; Philosophical Magazine Letters vol. 82 (2002) pp.543-551; Acta Materialia vol. 50 (2002) pp. 3845-3857).

The stacking faults are formed because a stacking order of a closelypacked side is changed unlike a normal stacking order, and it is knownthat they are mostly formed due to plastic deformation. It is difficultto form stacking faults if stacking fault energy is high, and thus,strain hardening required as a press material is not high. Accordingly,since pure aluminum or copper has high stacking fault energy, energysupplied during a room temperature process is mostly converted intoheat. Thus, it is difficult to accumulate internal deformations, and adriving force for nucleation is reduced during recrystallization.However, in the magnesium alloy of the present invention, the group IIIband IIIa elements are alloyed with magnesium acting as a matrix element,thereby reducing the stacking fault energy of the intermetallic compoundto provide ductility. Additionally, fine second phases promotenucleation during a reheating process to make fine grains. Intermetalliccompound particles prevent grain growth at a recrystallizationtemperature or higher.

Based on the above description, the present inventor came to aconclusion that when a group IIIa is alloyed with magnesium to form asolid solution having low stacking fault energy, when a group IIIb isadded to the solid solution to increase a solid-solution strengtheningeffect, and when a group IIb and other miniaturized elements are addedto form a structure which contains an intermetallic compound coherentthereto, it is possible to create a material having excellent strainhardening ability, fineness through recrystallization by heat treatment,and improved anisotropy.

Hereinafter, a detailed description will be given of elements andcompositions of the wrought magnesium alloy according to the presentinvention.

The group IIIa, that is, an essential element in the present invention,includes Sc, Y, lanthanides, and actinides. In this regard, it ispreferable that Sc, Y, or lanthanides be employed alone or incombination instead of actinides radiating radioactive rays. They aresolid-solved in Mg, thus reducing a c/a ratio to increase ductility andreducing the stacking fault energy to increase the driving force fornucleation by recrystallization. Furthermore, particles, which exist ina form of Mg₅RE at high temperatures during a solidification process,form prism-shaped plate particles having a HCP structure, that is, aDO₁₉ lattice structure, such as Mg₃RE or Mg₁₇RE₅, at about 550° C.through a peritectic transformation. (RE is an abbreviated form ofrare-earth elements belonging to the group if IIIa) Thereby, theparticles have a high reinforcement effect and are coherent to thematrix, and consequently, they do not act as the fracture initiationsite. After a rolling process, the particles may be compacted into arod, a sphere, or a cube.

In the present invention, a eutectic phase, which is not solid-solvedafter a diffusion heat-treatment, is finely dispersed during extrusionand rolling processes, thereby preventing grain growth duringheat-treatment and acting as a site for nucleation byre-crystallization. When the amount of the group IIIa is less than 0.1%,the second phase is formed in an insufficient amount. When the amount ismore than 1.5%, a fineness effect is saturated, and consequently,elongation is reduced and a production cost increases. This is thereason why that amount is limited.

The group IIIb includes B, Al, Ga, In, and Ti. Since Ga, In, and Tlhaving a low melting point form a low melting point eutectic phase, itis preferable to employ only Al or a mixture of B and Al. The group IIIbforms a fine deposit and thus contributes to reinforcement of thematrix. Al is used as a main alloy element. Since B has a low solidsolubility to magnesium and forms a high melting point compound, such asB₂Y, B₃Y₂, or B₅Y₃, it is employed in conjunction with Al in an amountof 0.010% or less so as to make the fine grains.

In the present invention, Al of the group IIIb is solid-solved in Mg toincrease corrosion resistance and to prevent the growth of a dendritemicrostructure, thereby making a cast microstructure fine. Furthermore,since Al forms fine cubes, such as Al RE or Al₃RE during thesolidification process and increases ductility of the matrixmicrostructure, it is possible to produce goods having high strength andexcellent ductility. When the amount of Al is less than 1.0%, it isdifficult to assure the desirable reinforcement effect. When the amountis more than 4.0%, since an unstable rod- or plate-shaped Al₂Mg₃ orAl₁₂Mg₁₇ phase is enlarged in a grain boundary, even though roomtemperature strength is high, high temperature strength and corrosionresistance are reduced. This is the reason why that amount is limited.

In order to make the fine grains and to help form the intermetalliccompound, 0.35% or less of group IIa, group IVa, group VIIa, or groupIVb is selectively employed alone or in combination, and 1.0% or less ofgroup IIb is employed alone or in combination.

The group IIa, group IVa, and group VIIa are used as a supplementalagent of the group IIIa and group IIIb. In the group IIa, it ispreferable to use Ca and Sr. Since Be, Ba, and Ra make toxic gases, theycan be used only if a special ventilation device is adopted. Ca and Srare particularly useful to make a fine cast structure in the casting abillet having a diameter of 200 mm or more in the present invention, andform disk-shaped particles, such as (Mg, Al)₂Ca, thereby improving areinforcement effect.

In the group IVa, Ti, Zr, and Hf are most frequently employed, and Rf isadded, using a protection device in unavoidable cases, because of theemission of radioactive rays. The group IVb makes a cast microstructurefine, and Si and Ge are most frequently employed because a melting pointis high and it is easy to handle. A grain fining effect depends on theamount of each element added. That is, Zr, Si, and Ca make the grainshave fine sizes of microns corresponding to the reciprocal of 52, 19,and 15 microns.

Mn of the group VIIa is a cheap alloy element, prevents the formation ofAl₁₂Mg₁₇ and Al₂Mg₃ phases, and promotes the formation of hightemperature cubic Al₂Y to contribute to the fining of the grains and theimprovement of corrosion resistance. Tc and Re of the group VIIa arecostly and thus are used in unavoidable cases.

The group IIa, group IVa, group VIIa, and group IVb elements have lowsolid solubility to magnesium, and thus, if they are excessively added,segregation occurs or coarse particles having high brittleness areformed when a cooling rate is low after a casting process. Accordingly,the amount is limited to 0.35% or less.

The group IIb includes Zn, Cd, and Hg. Since Hg is toxic to humans whenbreathing, use of Hg is limited, and it is used in conjunction with anadditional protection device. When Zn and Cd are added alone or incombination, a stacking fault structure is formed in a magnesium matrixmicrostructure to bring about strain hardening, and Zn and Cd aresmoothly solid-solved with the group IIIa and group IIIb elements topromote the formation of cubic particles, such as (Mg, Zn)₅RE, Zn₆Mg₂RE,or (Mg, Zn)₁₇RE₃. However, the excessive amount of Zn and Cd increasesgas solid solubility, thereby reducing corrosion resistance or platingworkability and brining about the occurrence of hot tear and gravityseparation phenomena. Hence, the amount is limited to 1.0% or less, andpreferably, 0.65% or less.

Hereinafter, a method of producing a plate using a magnesium alloy slabaccording to the present invention is described in detail through thefollowing example which is set forth to illustrate, but is not to beconstrued as the limit of the present invention.

(a) A magnesium raw material is melted, and an alloy or a master alloyis added to the molten magnesium in a mixed gas atmosphere of SF₆ and Aror CO₂, or an Ar gas atmosphere while being blocked from contact withatmospheric air. Generally, a slab for a magnesium alloy plate isproduced through mold casting, Direct Chilled casting, continuouscasting, or strip casting processes.

In the present example, a mold, in which a cavity having a thickness of30 mm, a width of 250 mm, and a height of 400 mm is formed, is preheatedin a heating furnace heated to about 200° C. A molten magnesium alloy ispoured into the mold at 710-760° C., and then machined so as to removesurface defects from a cast product.

(b) Diffusion annealing is conducted at 250-450° C. so that the durationtime is 1 min/mm or more with respect to the thickness of the slab. Whenthe heating temperature is less than 250° C. or the duration time isless than 1 min/mm, the inside of the slab is insufficiently heated,thus forming cracks on a surface or on an edge during a rolling process.It is preferable to heat the slab at 350-400° C. so as to reduce thediffusion time. When the heating temperature is more than 450° C., afree low melting point eutectic phase may be formed during the diffusionannealing. At this stage, the eutectic phase may be remelted and thusseparated from the slab. Accordingly, the molten eutectic phase maycling to a rolling roll. When the amount of the alloy element is large,the duration time and the heating temperature increase to improveworkability.

(c) Initial coarse rolling is conducted once or more in a reductionratio of 20% or less each time so as to fracture a coarse castmicrostructure of a material, which is subjected to diffusion annealing,and to remove fine segregation. After the completion of a rollingprocess, a process annealing is conducted once or more at 200-450° C. sothat the duration time is 1 min/mm or more with respect to the thicknessof a slab. When the heating temperature is less than 200° C. or theduration time is less than 1 min/mm, the inside of the slab isinsufficiently heated, thus forming cracks on a surface or on an edgeduring the rolling process. In the initial coarse rolling, when areduction ratio is more than 20%, cracks may be formed at a grainboundary of a cast microstructure. At this stage, a surface temperatureof a rolling roll must be maintained at 50-150° C. so as to prevent theformation of fine surface cracks caused by the quenched slab while theslab is in contact with the roll. When the temperature of the rollingroll is more than 150° C., delamination, in which a portion of a rollingmaterial clings to the rolling roll and then delaminates, occurs duringthe rolling process, thus roughening the surface of the slab. If theplate is not excessively cooled after the initial coarse rolling, it ispossible to conduct the rolling process again without reheating.

(d) When the cast microstructure of the slab is fractured, a secondrolling process is repeatedly conducted in a reduction ratio of 50% orless each time until the desired thickness is gained. At this stage, thereduction ratio depends on the capacity of a motor of a rolling mill, aheat emitting state of a plate during a reduction process, elasticdeformation of the rolling roll, and flatteness of the plate. It ispreferable that a second process annealing be repeatedly conducted at200-450° C. each time while a duration time is maintained at 1 min/mm ormore during the second rolling process. However, in the second rollingprocess, a rolled microstructure becomes fine, causing crack resistance.Furthermore, in some cases, it is possible to conduct cold rolling.Thus, annealing is not necessarily conducted every rolling process.

(e) After the final rolling process is completed, the final annealing isconducted at 180-350° C. while a duration time is maintained at 1 min/mmor more, which depends on the thickness, strength, and elongation of theplate. When the annealing temperature is high and the time is long,elongation increases but strength is reduced. Particularly, when theannealing temperature is more than 350° C., undesirably, yield strengthis significantly reduced.

Hereinafter, a detailed description will be given of extrusion using amagnesium alloy billet according to the present invention.

(a) A magnesium raw material is melted, and an alloy raw material or amaster alloy is added to the molten magnesium raw material in a mixedgas atmosphere of SF and Ar or CO₂, or an Ar gas atmosphere while beingblocked from contact with atmospheric air. Subsequently, a moltenmagnesium alloy is poured into a mold, having a diameter of 185 mm and alength of 650 mm, at 710-760° C. to form a billet, and is then processedso as to remove surface defects. Needless to say, it is possible toconduct a continuous casting in addition to a mold casting.

(b) Diffusion annealing is conducted at 250-450° C. while a durationtime is maintained at 1 min/mm or more with respect to the diameter ofthe billet so as to fracture a coarse cast microstructure of a castmaterial and to remove fine segregation. When a heating temperature isless than 250° C. or the duration time is less than 1 min/mm, stress isconcentrated on a grain boundary, and consequently, alligatoring mayoccur, cracking the material in a direction of the extrusion. It ispreferable to heat the material at 350-400° C. so as to reduce adiffusion time. When the heating temperature is more than 450° C., afree low melting point eutectic phase may be re-melted during thediffusion annealing and thus be separated from the material. When theamount of the alloy element is large, the duration time and the heatingtemperature increase to improve workability.

(c) The diffusion-annealed material is reheated in a heating furnace at250-400° C. to be extruded. An extruder has an extrusion speed of amaximum of 20 ml/min at an extrusion pressure of 850 MPa or more. If theextrusion is conducted at 500 MPa, the extrusion speed is significantlyreduced to 3-4 m/min. A temperature of a container is 300-450° C. Whenthe temperature is less than 300° C., many surface cracks are formed.When the temperature is more than 450° C., high temperature cracks ordeformations are significantly formed during the extrusion process. Thecontainer is heated at about 350° C., and an extrusion ratio istypically 10-100. Additionally, in the present invention, the materialmay be wound in a coil form during the extrusion process, and thus, itis possible to conduct reciprocating rolling.

(d) If the billet is very large or the cast microstructure is coarse, afirst extrusion is conducted to fracture the cast microstructure and todisperse a second phase, and a second extrusion is then conducted. Afterthe first extrusion, it is preferable to conduct a process annealing at200-450° C. while a duration time is maintained at 1 min/mm or more.However, during the first extrusion, the microstructure is made fine,causing crack resistance, and the reheating is implemented in thecontainer. Hence, annealing is not necessarily conducted.

(e) After the final rolling process is completed, if the material isrolled into a plate, the final annealing is conducted at 180-350° C.while a duration time is maintained at 1 min/mm or more, which dependson a thickness, strength, and elongation of the plate. When theannealing temperature is high and the time is long, elongation increasesbut strength is reduced. Particularly, when the annealing temperature ismore than 350° C., undesirably, yield strength is significantly reduced.Needless to say, when the plate and coil are annealed, the heattreatment may be implemented using a rapidly heating device, such as aheater, employing a gas nozzle, or an induction heater, instead of thefurnace. At this stage, since a heating rate is high, it is necessary toset the annealing temperature higher. In this regard, the annealingtemperature may deviate from the above range, without departing from thescope and concept of the invention.

As shown in following Table 1 and 2, wrought magnesium alloys of thepresent invention were rolled to obtain test results. They were testedafter being rolled into plates having a width of 150 mm and a thicknessof 1 mm.

Rectangular molds, which had a width of 80 mm, a length of 100 mm, and adepth of 45 mm, were formed, and edge cracks of the molds were observed,thereby achieving a forming test. Samples having an area of 80 mm×50 mmwere hung on a nylon thread as a hanger, and immersed in 200 cc of 2%HCl aqueous solution in a beaker. Thereby, gases, generated from thesamples, were dissolved in the solution. At this stage, weight reductionwas measured, thereby achieving evaluation of corrosion resistance. Theevaluation of formability is as follows. ◯ means that forming isachieved without cracks and local reduction of a thickness, Δ means thatcracks are not formed but a thickness deviation locally occurs, and xmeans that formability is very poor because of the formation of cracks.In the evaluation of characteristics of a wet plating process, ◯ means astate that plating thickness and adhesion of a plated surface areexcellent. Δ means a state that adhesion is fair, pinhole is notobserved, and plating thickness is ununiform. x means a state in whichthe pinholes are observed or plating layer comes off the surfacesomewhere in the specimen.

TABLE 1 Chemical component(at %)The balance includes impurities and MgProcess Results Group (dimension ∘: IIa, IVa, of mold, excellent GroupGroup VIIa, Group and unit is Mechanical Δ: fair No. IIIa IIIb IVb IIbmm) properties x: poor Note 1 Y 0.2 Al 2.00 Mn 0.13 Zn 0.40 30 × 250 ×¹T. 270 MPa ⁴F. ∘ ¹⁰I.S. Sc 0.11 400Mold ²Y. 224 MPa ⁵C.R. 3.5 casting³El 22% ⁶P. ∘ 2 Y 0.50 Al 1.50 Mn 0.15 Zn 0.50 Dia. T. 257 MPa ⁴F. ∘¹⁰I.S. Nd 0.01 Cd 0.10 185Billet Y. 213 MPa ⁵C.R. 3.2 casting El 20% ⁶P.∘ 3 Y 0.50 Al 3.00 Ca 0.01 Zn 0.30 Dia. T. 376 MPa ⁴F. ∘ ¹⁰I.S. Mn 0.10185Billet Y. 312 MPa ⁵C.R. 3.0 casting El 21% ⁶P. ∘ 4 Y 0.45 Al 2.50 Sr0.02 Zn 0.50 Dia. T. 355 MPa ⁴F. ∘ ¹⁰I.S. Sc 0.10 Mn 0.20 305Billet Y.287 MPa ⁵C.R. 3.3 Zr 0.08 casting El. 19% ⁶P. ∘ 5 Y 1.00 Al 3.00 Si 0.10Zn 0.40 Dia. T. 346 MPa ⁴F. ∘ ¹⁰I.S. La 0.05 Mn 0.15 Cd 0.10 185BilletY. 287 MPa ⁵C.R. 2.9 casting El. 18% ⁶P. ∘ 6 Y 0.40 Al 3.50 Zr 0.05 Zn0.50 Dia. T. 405 MPa ⁴F. ∘ ¹⁰I.S. Nd 0.03 185Billet Y. 336 MPa ⁵C.R. 3.4casting El. 18% ⁶P. ∘ 7 Y 0.30 Al 2.50 Zr 0.10 Zn 0.50 Dia. T. 331 MPa⁴F. ∘ ¹⁰I.S. B 0.005 185Billet Y. 275 MPa ⁵C.R. 3.2 casting El. 23% ⁶P.∘

TABLE 2 8 Y 0.15 Al 2.00 Ca 0.10 Zn 0.30 Dia. T. 245 MPa ⁴F. ∘ ¹⁰I.S. Mn0.10 185Billet Y. 203 MPa ⁵C.R. 3.8 casting El. 18% ⁶P. ∘ 9 Y 0.25 — Zr0.80 Zn 1.55 D360 × T. 285 MPa ⁴F. ∘ ¹¹C.S. t120D.C. Y. 253 MPa ⁵C.R.3.8 casting El. 16% ⁷failed Ni 10 Y 0.15 Al 0.90 — Zn 0.75 30 × 250 × T.261 MPa ⁴F. ∘ ¹¹C.S. 400Mold Y. 205 MPa ⁵C.R. 5.2 casting El. 18% layeroff 11 — Al 2.54 Mn 0.09 Zn 0.30 30 × 250 × T. 265 MPa ⁴F. x ¹¹C.S.400Mold Y. 185 MPa ⁵C.R. 4.1 (AZ31) casting El. 18% ⁹Pinhole ¹T: Tensilestrength ²Y: Yield point ³El.: Elongatin ⁴F.: formability, ⁵C.R.:corrosion weight loss rate, ⁶P.: plating, ⁷Failed Ni: impossible to forma Ni plating, ⁸Layer off: the plating layer comes off the surface,⁹Pinhole: formation of micro pinholes in the plating layer, ¹⁰I.S.:present inventive sample ¹¹C.S.: comparative sample (No. 1, 9, 10, 11 -rolling speed: 1.6 m/min, reduction ratio: 15% during a initial coarserolling and then 20-45%, No. 2~8 - rolling speed: 16-20 m/min, reductionratio: 30-67% after extrusion)

When evaluating corrosion resistance, weights of the beakers, in whichthe samples were contained, were measured every five minutes for 60 minusing a precision scale having an allowable margin of error of 1/1000 gto calculate a slope of weight reduction, thereby completing theevaluation of corrosion resistance. The higher slope brings aboutincreased weight reduction, resulting in poor corrosion resistance.

In Table 1, since a magnesium alloy of No. 11, which is producedaccording to a conventional method, has poor formability, cracks areformed during the forming process as shown in FIG. 3. In the wet platingprocess, activation treatment is conducted in liquid, an electroplatingprocess, such as a copper cyanide plating, a copper sulfate plating, ora nickel plating, is then implemented, and subsequently, a finalplating, such as a chromium plating or a precious metal plating, isimplemented. At this stage, if the pinholes are formed or a platinglayer comes off surface, the reliability of corrosion resistance issignificantly reduced.

INDUSTRIAL APPLICABILITY

As described above, in the present invention, a fine second phaseintermetallic compound is dispersed so as to significantly improve thepoor formability and corrosion resistance of a conventional magnesiumplate. Thereby, the magnesium plate has excellent properties as astructural material, and consequently, it is possible to apply themagnesium plate to structural materials used in portable electronicproducts, automobiles, or airplanes.

1. A wrought magnesium alloy having excellent formability and plating properties, which comprises 0.1-1.5 at % group IIIa 1.0-4.0 at % group IIIb, 0.35 at % or less of one selected from the group consisting of groups IIa, IVa, VIIa, IVb, and a mixture thereof, 1.0 at % or less of group IIb, and a balance of Mg and unavoidable impurities and thus has a second phase intermetallic compound.
 2. The wrought magnesium alloy as set forth in claim 1, wherein the group IIIa includes Sc, Y, and Lanthanides.
 3. The wrought magnesium alloy as set forth in claim 1, wherein the group IIIb includes Al and B.
 4. The wrought magnesium alloy as set forth in claim 1, wherein the group IIa includes Ca and Sr.
 5. The wrought magnesium alloy as set forth in claim 1, wherein the group IVa includes Ti, Zr, and Hf.
 6. The wrought magnesium alloy as set forth in claim 1, wherein the group VIIa includes Mn.
 7. The wrought magnesium alloy as set forth in claim 1, wherein the group IVb includes Si and Ge.
 8. The wrought magnesium alloy as set forth in claim 1, wherein the group IIb includes Zn and Cd, and a content of the group IIb is 0.65 at % or less.
 9. A method of producing a wrought magnesium alloy, comprising: extruding a magnesium alloy cast billet which comprises 0.1-1.5 at % group IIIa, 1.0-4.0 at % group IIIb, 0.35 at % or less of one selected from the group consisting of groups IIa, IVa, VIIa, IVb, and a mixture thereof 1.0 at % or less of group IIb, and a balance of Mg and unavoidable impurities and thus has a second phase intermetallic compound; and then rolling the magnesium alloy plate. 